Electronic junctions on edge

Two-dimensional materials such as graphene are attractive materials for making smaller transistors because they are inherently nanoscale and can carry high currents. However, graphene has no band gap and the transistors are “leaky”; that is, they are hard to turn off. Related transition metal dichalcogenides (TMDCs) such as molybdenum sulfide have band gaps. Transistors based on these materials can have high ratios of “on” to “off” currents. However, it is often difficult to make a good voltage-biased (p-n) junction between different TMDC materials. Li et al. succeeded in making p-n heterojunctions between two of these materials, molybdenum sulfide and tungsten selenide. They did this not by stacking the layers, which make a weak junction, but by growing molybdenum sulfide on the edge of a triangle of tungsten selenide with an atomically sharp boundary

Two-dimensional (2D) transition metal dichalcogenides (2D TMDCs) are of interest for electronics applications in that they offer tunabilty of several properties, including the band gap, band offset, carrier density, and polarity. (The bulk TMDCs have been known for a long time and have not evoked similar interest.) Heterostructures formed by vertical stacking of different 2D TMDCs have been realized via the transfer of their exfoliated or as-grown flakes (1, 2), where their properties are dominated by the stacking orientation and interlayer coupling strength. However, lateral heterostructures with edge contacts offer easier band offset tuning because the materials are more spatially separated. The direct growth of lateral heterojunctions is challenging because TMDC alloys are thermodynamically preferred (3). Recently, the MoS2-MoSe2 (4), WS2-WSe2 (4), WS2-MoS2 (2), and MoSe2-WSe2 (5) lateral heterostructures with interesting optical and electrical properties were obtained by one-pot synthetic processes. However, the interface regions for these lateral junctions are likely alloy structures because all of the precursors coexist in vapor phases during the growth. Such processes only allow the growth of heterostructures with either different metals or chalcogen, making it difficult to grow p-n heterostructures such as WSe2-MoS2.

We report the controlled epitaxial growth of WSe2-MoS2 lateral junction, where WSe2 is grown on substrates through van der Waals epitaxy, followed by the edge epitaxy of MoS2 along the W growth front. Two-step growth offers precise control to achieve the atomically sharp transition in compositions at the junction. Optical and microscopic characterizations revealed the detailed mechanisms for the regrowth (similar to living growth) for the 2D TMDC systems. The 2D lateral WSe2-MoS2 heterojunction was synthesized on c-plane sapphire substrates by sequential chemical vapor deposition (CVD) of WSe2 and MoS2 (Fig. 1A) (6, 7). To avoid the alloy reaction observed in one-pot synthesis, we first prepared single-crystalline triangular WSe2 monolayer requiring a higher growth temperature (925°C) and then performed the MoS2 growth at 755°C in a separate furnace. As we previously reported (7), the WSe2 growth proceeds from WSe2 seeds, followed by van der Waals epitaxy on sapphire. The crucial point for successful heterostructure synthesis without alloy formation is to control the relative vapor amount of MoO3 and S during the second-step MoS2 growth. The excess in Mo precursors enhances the MoS2 vertical growth, whereas the excess in S vapor promotes the formation of undesired WS2 at the interface (fig. S1).

(A) Schematic illustration of the sequential growth of the monolayer WSe2-MoS2 in-plane heterostructure. (B) As shown in the optical image, the WSe2 and MoS2 can be distinguished by their optical contrast. (C and D) High-resolution STEM images taken from the WSe2-MoS2 in-plane heterostructure. (E) Atomic model showing the interface structure between WSe2 and MoS2.

The morphology of in-plane heterostructures was examined by optical microscopy (OM) and photoluminescence (PL) and Raman spectroscopies. Figure 1B shows the OM images of the lateral WSe2-MoS2 heterojunctions. All of the WSe2 triangles are uniformly surrounded by MoS2, and the domain for WSe2 and MoS2 can be distinguished simply by their optical contrast. The lattice constant of WSe2 is 5.53% larger than MoS2 (5, 8), which might be one of the factors restricting the growth of MoS2 onto WSe2 basal planes. The Raman and PL spectra in fig. S2 verify the chemical composition of inner WSe2 and outer MoS2 and also reveals the formation of the seamless WSe2-MoS2 junction.

The annular dark field (ADF) image of the lateral WSe2-MoS2 junction obtained with scanning transmission electron microscopy (STEM) revealed that the ADF signal increased with the atomic number (Z) as ~Z1.7 (Fig. 1C) (9). Thus, the W (74), Mo (42), Se (34), and S (16) atoms could be distinguished by their intensity. The ADF image for another location (Fig. 1D) shows the atomic models corresponding to the obtained image. An atomically sharp interface between the WSe2-MoS2 junction was formed, where about 90% of W atoms are located at the interface bridging to two pair of Se atoms and one pair of S atoms, as depicted in Fig. 1E. In addition to the ADF image, the coherent interface is also identified by the electron energy loss spectroscopy (EELS) measurement. The EELS line scan (by monitoring the EELS spectra change across the heterojunction) shown in fig. S3 verified an atomically sharp change. These observations suggest that the growth starts from the replacement of Se atoms of WSe2 edge by S atoms.

Polarization-resolved second-harmonic generation (SHG) microscopy is sensitive to the crystal orientation and domain boundaries of surface layers (10–16). We used a back-reflection geometry with a linearly polarized pump laser (870 nm) normally incident on a triangular WSe2-MoS2 heterostructure sample and detected the SH intensities with polarizations parallel (IH) and perpendicular (IV) to the laser polarization (Fig. 2A) (see fig. S4 for the setup). The total SH intensity Itotal (sum of IV and IH) (Fig. 2C), which was generally uniform over the entire WSe2 and MoS2 domains, indicated that the surrounding MoS2 was also single crystalline without grain boundaries. The SH intensity also showed no suppression across the junction, which suggests that the MoS2 grew out from the edges of WSe2 without misorientation. The MoS2 regions showed slight variations in SH intensity that coincided well with the variations in PL intensity and peak energy (Fig. 3A, inset). As will be discussed below, we consider that the SH intensity variations in the surrounding MoS2 arise from the nonuniform strain distribution rather than presence of multiple grains.

(A) The intensity maps of the perpendicular IV and parallel IH components of the SH. The insert shows the optical image. The black double arrow line indicates the direction of incident laser polarization. (C and D) Maps of the total intensity Itotal and angle θ between the direction of laser polarization and the armchair direction of the sample, respectively. (B) The intensity maps of the IV and IH components of the SH for a multidomain WSe2-MoS2 junction. (E and F) Maps of the Itotal and θ of the area shown in (B). The scale bar for all is 5 μm.

(A) PL spectra of MoS2 and its corresponding spatial modulation as shown in the inset contour color map. (B) Spatial maps of the A1g and E2g peak position. (C) The extracted MoS2 Raman spectra from the WSe2-MoS2 junction, where the colored spectra corresponds to the inset contour color map in (A). (D and E) The maps showing PL intensity and the spatial modulation of photon energy for a selected heterojunction. (F) Spectra 1 to 9 were collected from the location marked in (D).

To gain more insight into the WSe2-MoS2 heterojunction growth, we calculated the crystal orientation using , where θ is the angle between the laser polarization direction and the nearest armchair axis of the sample (13). The map of θ (Fig. 2D) was uniform over the entire WSe2-MoS2 heterostructure, indicating that the outgrowing MoS2 was a single crystal with the same orientation as the inner WSe2. We have also performed SHG measurements on a WSe2-MoS2 heterojunction composed of multiple grains (Fig. 2, B and E). Although these grains had different orientations (θ map, Fig. 2F), the outgrowing MoS2 followed the orientation of the inner WSe2. These results strongly suggest that the outgrowth of MoS2 occurred through epitaxy off the edge of WSe2 and determined the crystal orientation, rather than the sapphire substrate. Fig. S5 demonstrates that the MoS2 monolayer also grew out from the prepatterned WSe2 monolayer.

We noticed that the MoS2 in WSe2-MoS2 heterostructures normally exhibits considerable PL energy differences at different locations (Fig. 3A, inset), where the contour color mapping shows the spatial distribution of PL energy ranging from 1.79 to 1.91 eV. The site with a higher PL energy always exhibited a higher intensity, as presented in the PL spectra (Fig. 3A). However, such a large variation in PL was not observed in isolated MoS2 monolayers occasionally found on the same sample (fig. S6), so the MoS2 PL variation is related to the heterostructure formation. We also performed Raman imaging for the MoS2 A1g and E2g frequencies (Fig. 3B). Compared with the inset of Fig. 3A, the location with a higher PL intensity and energy also exhibited higher A1g and E2g frequencies.

According to previous studies on the strain-dependent E2g and A1g Raman modes in MoS2 monolayer (17–20), we conclude that both the PL and Raman variations reflect the local strain distribution in the MoS2. The frequency upshift (downshift) of both Raman modes is associated with a compressive (tensile) strain. The spectra from the isolated MoS2 and the three corners of the triangular heterostructure have an identical PL energy of 1.86 eV, near that of 1.82 ± 0.02 eV from unstrained MoS2 (19). For simplicity, assuming the isolated MoS2 is nearly strain free, we could then map out the relative strain on MoS2; the tensile area is colored with green, cyan, and black, and the compressive area is colored with red and yellow in the inset of Fig. 3A. The representative Raman spectra under tensile and compressive strain (Fig. 3C) show that the MoS2 with the lowest PL energy (1.79 eV) was frequently observed from the area with a tensile strain. As the compressive strain increases, the PL peak shifted to a higher energy with a pronounced intensity increase. Symmetry breaking of the crystal (19, 20) broke the degeneracy in the MoS2 E2g Raman mode (subpeaks E′– and E′+) for the as-grown TMDC monolayer (black curve).

The strain variation likely originated from the lattice mismatch between MoS2 and WSe2. Figure 3A shows that the strain of MoS2 adjacent to WSe2 was tensile (cyan color) and then gradually changed to strain free (blue color), particularly at the corners. To balance the strain built upon the MoS2, some MoS2 areas exhibited compressive strain (red color). We estimated the strain (relative to the as-grown isolated MoS2) in the MoS2 region of the MoS2-WSe2 heterostructure to be 1.59 ± 0.25% for largest tensile strain and 1.1 ± 0.18% for largest compressive strain, based on the reported linear PL energy shift rate of 45 meV/% strain (17–22). Such a large strain difference induced by the lateral heterostructure could indicate the possibility of using monolayer TMDCs for straintronics (23).

In the inner triangular WSe2, the PL spectra showed a prominent direct band-gap emission at ~1.63 eV. By contrast to the outer MoS2 region, the PL energy and intensity in the WSe2 region did not show pronounced variations (Fig. 3, D and E). Additionally, the Raman frequencies of the WSe2 region were also relatively unchanged (fig. S7). Interestingly, Fig. 3D shows that the PL emission from the heterojunction interface was stronger and the PL enhancement was localized at the WSe2 side of the lateral interface. The PL spectra taken from the points marked as 1 through 9 in Fig. 3D are displayed in Fig. 3F, where two characteristic peaks—1.62 eV for WSe2 (points 1 to 3) and 1.85 eV for MoS2 (points 7 to 9)—were observed, respectively. The line scan was performed at a selected MoS2 area nearly free of strain. Noticeably, the PL spectrum for WSe2 adjacent to the heterojunction (point 3) showed a narrower and stronger peak at 1.62 eV. The WSe2-MoS2 vertical contact forms a type II band alignment (1, 24, 25), with an interband transition peak at ~1.59 eV. However, we did not detect any appreciable interband transition in PL measurements.

Our observation differs from that for the WS2-MoS2 heterojunction in a previous report (2), where a broader PL peak with an intermediate band gap energy, identified as the interband transition, was observed. Meanwhile, it was reported that the edges of WS2 monolayers exhibit extraordinary PL intensity. (26) Our STEM results showed that interfacial W is bonded to Se and S, respectively, from each side, and the interface structure is similar to the reported WS2 edges. We performed a separate study on gas phase sulfurization of isolated WSe2 monolayer triangles and found that the PL emission from WSe2 edge was largely enhanced after edge sulfurization (fig. S8). The enhancement of PL at the interface could be related to the chemical composition change in the junction.

To study the depletion region of the atomically sharp WSe2-MoS2 heterojunction, we used scanning Kelvin probe microscopy (SKPM) to directly extract the spatial distribution of the surface potentials. Figure 4A and its inset show the SKPM and atomic force microscopy (AFM) images for the junction, respectively. The color contrast in the SKPM image between the WSe2 and MoS2 regions revealed the distinct potential difference across the junction. SKPM allows the measurement of depletion width, but the actual built-in potential difference is not accurate because it is strongly affected by the surface adsorbates. The line profile in fig. S9 reveals that the junction depletion width is ~320 nm. The value agrees with that of the 100 to 500 nm estimated from the depletion approximation of solving Poisson’s equation based on the assumptions of no free carriers and constant dopant concentration in the depletion region (see the supplementary materials).

To investigate the electrical properties, the as-grown WSe2-MoS2 heterojunction was transferred onto a SiO2/Si substrate, and two contact metals, Pd and Ti/Au, were deposited on WSe2 and MoS2, respectively (see the supplementary materials for details). Figure 4B shows the OM image of the WSe2-MoS2 heterojunction, and Fig. 4C shows the current-voltage (I-V) curves of the heterojunction without (black) and with (red) white light illumination. The characteristic curve exhibits good rectification character, with a threshold voltage at about 0.9 V under forward bias (fig. S10). A photovoltaic effect with an open-circuit voltage Voc of 0.22 V and short-circuit current Isc of 7.7 pA under white light illumination (power density Ew of 1 mW/cm2) is shown in the inset of Fig. 4C. The nearly symmetric I-V curves and barely photovoltaic effect for individual WSe2 (contacted with Pd) and MoS2 (contacted with Ti/Au) in Fig. 4D corroborate that the p-n junction from the heterostructure is predominant, rather than the small Schottky barriers between metal and TMDCs.

We calculated the power conversion efficiency (PCE) of the device with the photon-to-electron conversion equation, PCE = ISCVOCFF/EWAC, where FF is the fill factor and AC is the effective area with energy conversion. The FF of 0.39 was extracted from the inset of Fig. 4C. The small FF might result from the high equivalent series resistance of the intrinsic TMDC layers. We estimated the maximum AC by considering both the depletion area of the junction and the adjacent diffusion area of each TMDC layers, giving rise to a maximum area about 32 μm2 (see the supplementary materials for details). The calculated PCE is at least 0.2%, comparable with the few-layer MoS2 vertical p-n junction (27) and monolayer lateral WSe2 p-n junction (28). The p-n junction of WSe2-MoS2 is further corroborated with the results from another device (fig. S11), where the carrier transport is clearly through the heterojunction interface.

Acknowledgments: L.-J.L. acknowledges support from King Abdullah University of Science and Technology (Saudi Arabia), Ministry of Science and Technology (MOST) and Taiwan Consortium of Emergent Crystalline Materials (TCECM), Academia Sinica (Taiwan), and AOARD-134137 (USA). Y.-C.L. and K.S. acknowledge support from the Japan Science and Technology Agency research acceleration program. W.-H.C. acknowledges the support from TCECM, MOST of Taiwan under grant NSC102-2119-M-009-002-MY3 and the Center for Interdisciplinary Science of Nation Chiao Tung University.